KR102099769B1 - High strength steel plate and high strength galvanized steel plate - Google Patents

Abstract

The high-strength steel sheet has a specific chemical composition, and in the range of 1/8 thickness to 3/8 thickness centered on 1/4 thickness from the surface of the plate thickness, by volume fraction, ferrite: 85% or less, bainite: 3 % Or more and 95% or less, tempering martensite: 1% or more and 80% or less, residual austenite: 1% or more and 25% or less, pearlite and coarse cementite: 5% or less in total, and fresh martensite: 5% or less, Tempered martensite or fresh marten among the entire grain boundaries of the retained austenite particles having a microstructure expressed, a solid carbon content in the residual austenite of 0.70 to 1.30 mass%, an aspect ratio of 2.50 or less, and a diameter per circle of 0.80 µm or more. The interface occupied by the site occupies less than 75%.

Description

High strength steel plate and high strength galvanized steel plate

The present invention relates to a high-strength steel sheet and a high-strength galvanized steel sheet.

In recent years, in high-strength steel sheets used in automobiles and the like, demands for further improving the impact resistance properties have been increasing. In addition, in high-strength steel sheets used in automobiles and the like, in order to obtain a complicated member shape, moldability such as ductility and hole expandability is also required.

For example, a high-strength cold-rolled steel sheet aimed at improving ductility and elongation flangeability is described in Patent Document 1, and a high-strength steel sheet aimed at improving toughness and HAZ toughness is described in Patent Document 2, and the shape freeze and A high-strength steel sheet aimed at improving workability is described in Patent Document 3. In addition, Patent Document 4 describes a high-strength hot-dip galvanized steel sheet for the purpose of improving sinter hardenability while securing ductility, and Patent Document 5 provides a high-strength hot-dip galvanized steel sheet for the purpose of improving mechanical cutting properties while securing ductility. Patent document 6 describes a high-strength hot-dip galvanized steel sheet for the purpose of improving workability.

However, in the conventional high-strength steel sheet, excellent moldability and impact resistance properties that are recently required are not compatible.

Japanese Patent No. 5463685 Japanese Patent Publication No. 2014-9387 International Publication No. 2013/018741 International Publication No. 2013/047821 International Publication No. 2013/047739 Japanese Patent Publication No. 2009-209451

An object of the present invention is to provide a high-strength steel sheet and a high-strength galvanized steel sheet having excellent moldability and impact resistance.

The present inventors have made extensive studies to solve the above problems. As a result, it was found that it is important not only to optimize the chemical composition and the volume fraction of the microstructure, but also to arrange the coarse residual austenite, which is the starting point of destruction, to be as close to the tempering martensitic and fresh martensitic as possible. It has also been found that the control of the localization of Mn in the production process is very important for the batch control of such retained austenite, tempering martensite and fresh martensite.

In general, in order to control the microstructure volume fraction of a high-strength steel sheet, the mother phase is composed of austenite particles during residence near the highest heating temperature of annealing after cold rolling, and the subsequent cooling conditions and the like are adjusted. In the region where Mn is unevenly distributed, the austenite particles are coarsened during the residence, so that a structure in which coarse residual austenite and fresh martensite are adjacent to each other and mixed during cooling is obtained. In tempering after annealing, almost all of the fresh martensite becomes tempered martensite, but since the arrangement of the tissue is not changed in tempering, in the microstructure after tempering, coarse residual austenite and tempering martensite or fresh martensite They are adjacent and mixed. For example, coarse residual austenite is present to be surrounded by tempering martensite. In a high-strength steel sheet having such a microstructure, fracture tends to occur starting from the interface between coarse residual austenite and tempering martensite or fresh martensite.

Conventionally, for the optimization of the chemical composition and the volume fraction of the microstructure, conditions for annealing and tempering have been proposed, but since the localization of Mn proceeds with phase transformation at a relatively high temperature, only adjustment of these conditions is necessary. Mn does not control localization. The inventors of the present application conducted extensive investigations to suppress the localization of Mn, and as a result, can suppress the localization of Mn in the cooling process of hot rolling and the heating process of annealing, and suppress the localization of Mn to close to the highest heating temperature. It has been found that during the residence in Esau, the mother phase can be composed of fine and homogeneously dispersed austenite particles. By constituting fine and homogeneous dispersed austenite particles, after phase transformation with cooling, residual austenite and tempered martensite or fresh martensite are separated from each other by bainite and ferrite, etc. It becomes difficult to adjoin the tempering martensite or fresh martensite. Even if residual austenite surrounded by tempered martensite is present, the residual austenite is fine, making it difficult to be a starting point for destruction. Based on these findings, the present inventors considered all aspects of the invention shown below.

(One)

In mass%,

C: 0.075 to 0.400%,

Si: 0.01 to 2.50%,

Mn: 0.50 to 3.50%,

P: 0.1000% or less,

S: 0.0100% or less,

Al: 2.000% or less,

N: 0.0100% or less,

O: 0.0100% or less,

Ti: 0.000 to 0.200%,

Nb: 0.000 to 0.100%,

V: 0.000 to 0.500%,

Cr: 0.00 to 2.00%,

Ni: 0.00 to 2.00%,

Cu: 0.00 to 2.00%,

Mo: 0.00 to 1.00%,

B: 0.0000 to 0.0100%,

W: 0.00 to 2.00%,

Ca, Ce, Mg, Zr, La and one or more selected from the group consisting of REM: 0.0000 to 0.0100% in total,

Balance: Fe and impurities, also

Parameter Q0 represented by (Equation 1): 0.35 or more

Has a chemical composition represented by,

In the range of 1/8 thickness to 3/8 thickness centering on 1/4 thickness from the surface of the plate thickness, in a volume fraction,

Ferrite: 85% or less,

Bainite: 3% or more and 95% or less,

Tempering martensite: 1% or more and 80% or less,

Residual austenite: 1% or more and 25% or less,

Pearlite and coarse cementite: 5% or less in total, and

Fresh martensite: 5% or less

Has a micro-organism represented by,

The dissolved carbon content in the retained austenite is 0.70 to 1.30% by mass,

A high-strength steel sheet having an aspect ratio of 2.50 or less, and a ratio of the interface between the tempered martensite or fresh martensite and 75% or less of the total grain boundaries of the retained austenite particles having a circle equivalent diameter of 0.80 µm or more.

Q0 = Si + 0.1Mn + 0.6Al ... (Equation 1)

(In (Equation 1), Si, Mn and Al are the content of each element in mass%)

(2)

In mass%,

Ti: 0.001 to 0.200%,

Nb: 0.001 to 0.100% and

V: 0.001 to 0.500%

The high-strength steel sheet according to (1), characterized by containing one or two or more selected from the group consisting of.

(3)

In mass%,

Cr: 0.01 to 2.00%,

Ni: 0.01 to 2.00%,

Cu: 0.01 to 2.00%,

Mo: 0.01 to 1.00%,

B: 0.0001 to 0.0100% and

W: 0.01 to 2.00%

The high-strength steel sheet according to (1) or (2), characterized in that it contains one or two or more selected from the group consisting of.

(4)

In mass%,

The high-strength steel sheet according to any one of (1) to (3), characterized in that it contains 0.0001 to 0.0100% of one or more selected from the group consisting of Ca, Ce, Mg, Zr, La and REM.

(5)

The high-strength steel sheet according to any one of (1) to (4), characterized in that the density of residual austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more is 5.0 × 10 ¹⁰ particles / m 2 or less.

(6)

A high-strength galvanized steel sheet characterized in that a galvanized layer is formed on the surface of the high-strength steel sheet according to any one of (1) to (5).

(7)

The high-strength galvanized steel sheet according to (6), wherein the Fe content in the galvanized layer is 3.0% by mass or less.

(8)

The high-strength galvanized steel sheet according to (6), wherein the Fe content in the galvanized layer is 7.0% by mass or more and 13.0% by mass or less.

According to the present invention, since the relationship between the retained austenite and tempering martensite and fresh martensite is appropriate, excellent moldability and impact resistance can be obtained.

Hereinafter, embodiments of the present invention will be described.

(First embodiment)

First, the chemical composition of the high-strength steel sheet according to the first embodiment of the present invention will be described. Although the details will be described later, the high-strength steel sheet according to the first embodiment is produced through a hot rolling process, a pickling process, a cold rolling process, an annealing process, a bainite transformation process, a martensite transformation process, and a tempering process. Therefore, the chemical composition of the high-strength steel sheet is suitable for these treatments as well as the properties of the high-strength steel sheet. In the following description, "%" which is a content unit of each element contained in the high-strength steel sheet means "mass%" unless otherwise specified. High-strength steel sheet according to the present embodiment, C: 0.075 to 0.400%, Si: 0.01 to 2.50%, Mn: 0.50 to 3.50%, P: 0.1000% or less, S: 0.0100% or less, Al: 2.000% or less, N: 0.0100% or less, O: 0.0100% or less, Ti: 0.000 to 0.200%, Nb: 0.000 to 0.100%, V: 0.000 to 0.500%, Cr: 0.00 to 2.00%, Ni: 0.00 to 2.00%, Cu: 0.00 to 2.00 %, Mo: 0.00 to 1.00%, B: 0.0000 to 0.0100%, W: 0.00 to 2.00%, one selected from the group consisting of Ca, Ce, Mg, Zr, La and rare earth metal (REM) or 2 or more types: 0.0000 to 0.0100% in total, the balance: Fe and impurities, and also has a chemical composition represented by parameter Q0: 0.35 or more, represented by (Formula 1). As an impurity, what is contained in raw materials, such as ore and scrap, and what is contained in a manufacturing process are illustrated.

Q0 = Si + 0.1Mn + 0.6Al ... (Equation 1)

(In (Formula 1), Si, Mn and Al are the contents of each element in mass%)

(C: 0.075 to 0.400%)

C stabilizes austenite and obtains retained austenite, thereby increasing strength and formability. When the C content is less than 0.075%, residual austenite is not obtained, and it is difficult to ensure sufficient strength and formability. Therefore, the C content is 0.075% or more. In order to obtain more excellent strength and formability, the C content is preferably 0.090% or more, and more preferably 0.100% or more. On the other hand, if the C content is more than 0.400%, spot weldability deteriorates remarkably. Therefore, the C content is 0.400% or less. In order to obtain good spot weldability, the C content is preferably 0.320% or less, and more preferably 0.250% or less.

(Si: 0.01 to 2.50%)

Si stabilizes residual austenite by suppressing the production of iron-based carbides in the steel sheet, thereby increasing strength and formability. When the Si content is less than 0.01%, in a bainite transformation step, coarse iron-based carbides are generated in large quantities, and strength and formability deteriorate. Therefore, the Si content is 0.01% or more. In order to obtain superior strength and formability, the Si content is preferably 0.10% or more, and more preferably 0.25% or more. On the other hand, Si embrittles the steel or lowers the impact resistance properties of the steel sheet. When the Si content is more than 2.50%, embrittlement is remarkable, and trouble such as cracking of the cast slab is likely to occur. Therefore, the Si content is 2.50% or less. In order to obtain good impact resistance properties, the Si content is preferably 2.25% or less, and more preferably 2.00% or less.

(Mn: 0.50 to 3.50%)

Mn increases the quenchability of the steel sheet to increase the strength. If the Mn content is less than 0.50%, it is difficult to ensure a sufficiently high tensile maximum strength because a large amount of soft tissue is formed during cooling after annealing. Therefore, the Mn content is 0.50% or more. In order to obtain higher strength, the Mn content is preferably 0.80% or more, and more preferably 1.00% or more. On the other hand, Mn embrittles the steel or deteriorates the spot weldability. When the Mn content is more than 3.50%, a coarse Mn thickening portion is generated in the central portion of the sheet thickness of the steel sheet, embrittlement is likely to occur, and trouble such as cracking of the cast slab is likely to occur. Therefore, the Mn content is 3.50% or less. In order to obtain good spot weldability, the Mn content is preferably 3.20% or less, and more preferably 3.00% or less.

(P: 0.1000% or less)

P is not an essential element, and is contained as an impurity in steel, for example. Since P embrittles the steel or embrittles the melt generated by spot welding, the lower the P content, the better. When the P content is more than 0.1000%, embrittlement is remarkable, and trouble such as cracking of the slab is likely to occur. Therefore, the P content is 0.1000% or less. In order to suppress the embrittlement of the melted portion and obtain excellent weld joint strength, the P content is preferably 0.0400% or less, and more preferably 0.0200% or less. It is expensive to reduce the P content, and if it is desired to reduce it to less than 0.0001%, the cost is remarkably increased. For this reason, the P content may be 0.0001% or more, and preferably 0.0010% or more from the viewpoint of cost.

(S: 0.0100% or less)

S is not an essential element and is contained as an impurity in steel, for example. S is low when S content is low in order to form coarse MnS in conjunction with Mn to reduce formability such as ductility, hole expandability, stretch flangeability and bendability, or deteriorate spot weldability. The better. When the S content is more than 0.0100%, the deterioration of formability is remarkable. Therefore, the S content is 0.0100% or less. In order to obtain good spot weldability, the S content is preferably 0.0070% or less, and more preferably 0.0050% or less. It is expensive to reduce the S content, and if it is desired to reduce it to less than 0.0001%, the cost is significantly increased. For this reason, the S content may be 0.0001% or more, and the cost is preferably 0.0003% or more, and more preferably 0.0006% or more.

(Al: 2.000% or less)

In order for Al to embrittle the steel or deteriorate the spot weldability, the lower the Al content, the better. When the Al content is more than 2.000%, embrittlement is remarkable, and troubles such as cracking of the slab are likely to occur. Therefore, the Al content is 2.000% or less. In order to obtain good spot weldability, the Al content is preferably 1.500% or less, and more preferably 1.300% or less. It is expensive to reduce the Al content, and if it is desired to reduce it to less than 0.001%, the cost rises significantly. For this reason, the Al content may be 0.001% or more. Al is effective as a deoxidizing material, and in order to sufficiently obtain the effect of deoxidation, the Al content is preferably 0.010% or more. Al may be included for the purpose of stabilizing residual austenite in order to suppress the formation of coarse carbides. For stabilization of the retained austenite, the Al content is preferably 0.100% or more, and more preferably 0.250% or more.

(N: 0.0100% or less)

N is not an essential element and is contained as an impurity in steel, for example. The lower the N content is, the lower the N content is, in order to form coarse nitrides to reduce formability such as ductility, hole expandability, stretch flangeability, and bendability, or to generate blow holes during welding. . When the N content is more than 0.0100%, deterioration of moldability is remarkable. Therefore, the N content is 0.0100% or less. In order to suppress blow holes more reliably, the N content is preferably 0.0075% or less, and more preferably 0.0060% or less. It is expensive to reduce the N content, and if it is desired to reduce it to less than 0.0001%, the cost is remarkably increased. For this reason, the N content may be 0.0001% or more, and the cost is preferably 0.0003% or more, and more preferably 0.0005% or more.

(O: 0.0100% or less)

O is not an essential element and is contained as an impurity in steel, for example. Since O forms an oxide and deteriorates formability such as ductility, hole expandability, stretch flangeability and bendability, the lower the O content, the better. When the O content is more than 0.0100%, deterioration of moldability is remarkable. Therefore, the O content is 0.0100% or less, preferably 0.0050% or less, and more preferably 0.0030% or less. It is expensive to reduce the O content, and if it is desired to reduce it to less than 0.0001%, the cost is significantly increased. For this reason, the O content may be 0.0001% or more.

(Parameter Q0: 0.35 or more)

Although the details will be described later, in the tempering step after the martensitic transformation step, there is a fear that residual austenite decomposes to bainite, pearlite, or coarse cementite during the heat treatment. Si, Mn and Al are particularly important elements to suppress the decomposition of retained austenite and to increase moldability, and if the parameter Q0 represented by (Formula 1) is less than 0.35, the above effect is not obtained. Therefore, the parameter Q0 is 0.35 or more, preferably 0.60 or more, and more preferably 0.80 or more.

Q0 = Si + 0.1Mn + 0.6Al ... (Equation 1)

(In (Formula 1), Si, Mn and Al are the contents of each element in mass%)

Ti, Nb, V, Cr, Ni, Cu, Mo, B, W, Ca, Ce, Mg, Zr, La, and REM are not essential elements and may be appropriately contained in a high-strength steel sheet in a predetermined amount as a limit. Element.

(Ti: 0.000 to 0.200%)

Ti contributes to an increase in the strength of the steel sheet by strengthening precipitation, strengthening fine particles by suppressing the growth of ferrite grains, and enhancing potential through suppressing recrystallization. Although the desired purpose is achieved even when Ti is not included, the Ti content is preferably 0.001% or more, and more preferably 0.010% or more, in order to sufficiently obtain these effects. However, when the Ti content is more than 0.200%, Ti carbonitride precipitates excessively, and the moldability may deteriorate. For this reason, the Ti content is 0.200% or less. From the viewpoint of formability, the Ti content is preferably 0.120% or less.

(Nb: 0.000 to 0.100%)

Nb contributes to an increase in strength of the steel sheet by strengthening precipitation, strengthening fine particles by suppressing the growth of ferrite grains, and enhancing potential through suppressing recrystallization. Although the desired purpose is achieved even when Nb is not included, in order to sufficiently obtain these effects, the Nb content is preferably 0.001% or more, and more preferably 0.005% or more. However, when the Nb content is more than 0.100%, the Nb carbonitride precipitates excessively and the moldability may deteriorate. For this reason, the Nb content is 0.100% or less. From the viewpoint of formability, the content of Nb is preferably 0.060% or less.

(V: 0.000 to 0.500%)

V contributes to an increase in the strength of the steel sheet by strengthening precipitation, strengthening fine particles by suppressing the growth of ferrite grains, and enhancing potential through suppressing recrystallization. Although the desired purpose is achieved even when V is not included, the V content is preferably 0.001% or more, and more preferably 0.010% or more, in order to sufficiently obtain these effects. However, when the V content is more than 0.500%, V carbonitride precipitates excessively and the moldability may deteriorate. For this reason, the V content is 0.500% or less, and more preferably 0.350% or less.

(Cr: 0.00 to 2.00%)

Cr improves quenchability and is effective for high strength. Although the intended purpose is achieved even when Cr is not included, the Cr content is preferably 0.01% or more, and more preferably 0.10% or more, in order to sufficiently obtain these effects. However, when the Cr content is more than 2.00%, workability in hot may be impaired, resulting in a decrease in productivity. For this reason, the Cr content is 2.00% or less, and preferably 1.20% or less.

(Ni: 0.00 to 2.00%)

Ni suppresses phase transformation at high temperature, and is effective for high strength. Although the intended purpose is achieved even when Ni is not included, the Ni content is preferably 0.01% or more, and more preferably 0.10% or more, in order to sufficiently obtain these effects. However, if Ni content exceeds 2.00%, weldability may be impaired. For this reason, the Ni content is 2.00% or less, and preferably 1.20% or less.

(Cu: 0.00 to 2.00%)

Cu is a fine particle and is present in steel to increase strength. Although the intended purpose is achieved even when Cu is not included, in order to sufficiently obtain these effects, the Cu content is preferably 0.01% or more, and more preferably 0.10% or more. However, if the Cu content is more than 2.00%, weldability may be impaired. For this reason, the Cu content is 2.00% or less, and preferably 1.20% or less.

(Mo: 0.00 to 1.00%)

Mo suppresses phase transformation at high temperature, and is effective for high strength. Although the intended purpose is achieved even when Mo is not included, in order to sufficiently obtain these effects, the Mo content is preferably 0.01% or more, and more preferably 0.05% or more. However, when the Mo content is more than 1.00%, the workability during hot may be impaired, resulting in a decrease in productivity. For this reason, the Mo content is 1.00% or less, and preferably 0.50% or less.

(B: 0.0000 to 0.0100%)

B suppresses phase transformation at high temperature and is effective for high strength. Although the desired purpose is achieved even if B is not included, in order to sufficiently obtain these effects, the B content is preferably 0.0001% or more, and more preferably 0.0005% or more. However, if the B content is more than 0.0100%, workability in hot may be impaired, resulting in a decrease in productivity. For this reason, the B content is 0.0100% or less, and more preferably, the B content is 0.0050% or less.

(W: 0.00 to 2.00%)

W suppresses phase transformation at a high temperature and is effective for high strength. Although the desired purpose is achieved even when W is not included, the W content is preferably 0.01% or more, and more preferably 0.10% or more, in order to sufficiently obtain these effects. However, when the W content is more than 2.00%, workability in hot may be impaired and productivity may decrease. For this reason, W content is 2.00% or less, preferably 1.20% or less.

(Ca, Ce, Mg, Zr, La and one or more selected from the group consisting of REM: 0.0000 to 0.0100% in total)

REM refers to an element belonging to the lanthanoid family. For example, REM or Ce is added as a misch metal, and may contain lanthanoid-based elements in a complex other than La or Ce. Even if a lanthanoid-based element other than La or Ce is included, the effect of the present invention is exhibited. Even if metal La or Ce is included, the effect of the present invention is exhibited.

Ca, Ce, Mg, Zr, La and REM are effective for improving moldability. Although the desired purpose is achieved even if Ca, Ce, Mg, Zr, La and REM are not included, in order to sufficiently obtain these effects, the total content of Ca, Ce, Mg, Zr, La and REM is preferably 0.0001% It is the above, More preferably, it is 0.0010% or more. However, if the total content of Ca, Ce, Mg, Zr, La and REM is more than 0.0100%, there is a risk of impairing ductility. For this reason, the total content of Ca, Ce, Mg, Zr, La and REM is 0.0100% or less, preferably 0.0070% or less.

As impurities, H, Na, Cl, Sc, Co, Zn, Ga, Ge, As, Se, Y, Zr, Tc, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Te, Cs, Hf , Ta, Re, Os, Ir, Pt, Au, and Pb in a total amount of 0.0100% or less is acceptable.

Next, the microstructure of the high-strength steel sheet according to the first embodiment will be described. The high-strength steel sheet according to the present embodiment has a volume fraction of ferrite: 85% or less, bainite: 3, in a range of 1/8 to 3/8 thickness centered on 1/4 thickness from the surface of the plate thickness. % Or more and 95% or less, tempered martensite: 1% or more and 80% or less, residual austenite: 1% or more and 25% or less, pearlite and coarse cementite: 5% or less in total, and fresh martensite: 5% or less Has a micro-structure.

(Ferrite: 85% or less)

Ferrite has excellent ductility. However, since the strength of ferrite is low, when the volume fraction of ferrite is more than 85%, sufficient tensile maximum strength cannot be obtained. For this reason, the volume fraction of ferrite is 85% or less. In order to obtain a higher tensile maximum strength, the volume fraction of ferrite is preferably 75% or less, more preferably 65% or less. Although the desired purpose is achieved even when ferrite is not included, in order to obtain good ductility, the volume fraction of ferrite is preferably 5% or more, and more preferably 10% or more.

(Bainite: 3% or more and 95% or less)

Bainite is a structure having an excellent balance of strength and formability. If the volume fraction of bainite is less than 3%, a good balance of strength and formability cannot be obtained. Therefore, the volume fraction of bainite is 3% or more. Since the volume fraction of retained austenite improves with the formation of bainite, the volume fraction of bainite is preferably 7% or more, and more preferably 10% or more. On the other hand, if the volume fraction of bainite is more than 95%, it becomes difficult to secure both tempering martensite and retained austenite, and a good balance of strength and formability cannot be obtained. For this reason, the volume fraction of bainite is 95% or less, and in order to obtain a balance of superior strength and formability, the volume fraction of bainite is preferably 85% or less, and more preferably 75% or less.

In addition, the bainite in the present invention is made of lath-shaped body-centered cubic (bcc) crystals and does not contain iron-based bainitic ferrite, fine bcc crystals, and granules made of coarse iron-based carbides. Ru bainite, upper bainite made of lath-shaped bcc crystals and coarse iron-based carbides, and lower bainite made of plate-like bcc crystals and fine iron-based carbides arranged in parallel therein.

(Tampering martensite: 1% to 80%)

Tempering martensite does not impair the impact resistance properties and greatly improves the tensile strength of the steel sheet. If the volume fraction of the tempered martensite is less than 1%, sufficient tensile strength cannot be obtained. Therefore, the volume fraction of tempering martensite is 1% or more. To obtain higher tensile strength, the volume fraction of tempering martensite is preferably 5% or more, and more preferably 10% or more. On the other hand, if the volume fraction of the tempered martensite is more than 80%, the interface between the tempered martensite and residual austenite is excessively increased, and the impact resistance characteristics are deteriorated. Therefore, the volume fraction of tempering martensite is 80% or less. In order to obtain better impact resistance properties, the volume fraction of the tempered martensite is preferably 73% or less, more preferably 65% or less.

(Residual austenite: 1% or more and 25% or less)

Residual austenite increases the balance between strength and ductility. If the volume fraction of retained austenite is less than 1%, a good balance of strength and ductility cannot be obtained. Therefore, the volume fraction of retained austenite is 1% or more. In order to obtain good moldability, the volume fraction of retained austenite is preferably 2.5% or more, and more preferably 4% or more. On the other hand, in order to make the volume fraction of the retained austenite exceed 25%, a degree of weldability of C is required. For this reason, the volume fraction of retained austenite is 25% or less. Residual austenite is transformed into hard martensite by an impact, and acts as a starting point for destruction. When the volume fraction of retained austenite exceeds 21%, martensite transformation is likely to occur. Therefore, the volume fraction of retained austenite is preferably 21% or less, and more preferably 17% or less.

The higher the solid solution carbon content in the retained austenite, the higher the stability of the retained austenite, and excellent impact resistance properties are obtained. If the amount of solid solution carbon in the retained austenite is less than 0.70% by mass, this effect cannot be sufficiently obtained. Therefore, the amount of solid solution carbon in the retained austenite is 0.70% by mass or more. The amount of solid solution carbon in the retained austenite is preferably 0.77% by mass or more, and more preferably 0.84% by mass or more. On the other hand, if the amount of solid solution carbon in the retained austenite is excessively high, transformation from residual austenite to martensite due to tensile deformation does not proceed sufficiently, and the work hardenability may decrease. If the solid carbon content in the retained austenite is more than 1.30% by mass, sufficient work hardenability cannot be obtained. Therefore, the amount of solid solution carbon in the retained austenite is 1.30% by mass or less, preferably 1.20% by mass or less, and more preferably 1.10% by mass or less.

Residual austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more are easily transformed into hard martensite by being subjected to impact, and are likely to act as a starting point for destruction. In particular, in the vicinity of the interface of the above-described residual austenite particles and the tempered martensite or fresh martensite, the retained austenite is constrained by the hard tempered martensite or fresh martensite, resulting in deformation. High deformation occurs on the night side, and the retained austenite is easily transformed into martensite. For this reason, the interface between the retained austenite and tempering martensite or fresh martensite is peeled off, and breakage is likely to occur.

And, among the total grain boundaries of the retained austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more, the proportion occupied by the interface with the tempered martensite or fresh martensite, that is, the portion in contact with the tempered martensite or fresh martensite When the ratio of, is more than 75%, fracture accompanying fracture of the interface is remarkable. Therefore, in this embodiment, this ratio is 75% or less. In this embodiment, since this ratio is 75% or less, the fracture starting from the residual austenite particles can be suppressed, and the impact resistance characteristics can be improved. In order to obtain better impact resistance properties, this ratio is preferably 60% or less, and more preferably 40% or less.

By reducing the coarse residual austenite particles that are the starting point for destruction, the impact resistance properties are further improved. From this viewpoint, the density of the retained austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more is preferably 5.0 × 10 ¹⁰ particles / m 2 or less, more preferably 3.0 × 10 ¹⁰ pieces / m 2 Is below.

(Fresh martensite: 5% or less)

Fresh martensite greatly improves the tensile strength, but, on the other hand, serves as a starting point for deterioration and deteriorates the impact resistance properties. When the volume fraction of fresh martensite is more than 5%, deterioration of impact resistance characteristics is remarkable. For this reason, the volume fraction of fresh martensite is 5% or less. In order to obtain excellent impact resistance properties, the volume fraction of fresh martensite is preferably 1% or less, and more preferably 0%.

(Perlite and coarse cementite: 5% or less in total)

Pearlite and coarse cementite deteriorate ductility. When the volume fraction of pearlite and coarse cementite exceeds 5% in total, deterioration in ductility is remarkable. For this reason, the volume fraction of pearlite and coarse cementite is 5% or less in total. Here, in the present embodiment, coarse cementite means cementite having a circle equivalent diameter of 1.0 µm or more. By observation under an electron microscope, the diameter of the equivalent of cementite can be easily measured, and it can be easily determined whether or not the cementite is coarse.

The volume fraction of ferrite, bainite, tempered martensite, fresh martensite, pearlite and coarse cementite can be measured using the method shown below. Samples were taken with a plate thickness section parallel to the rolling direction of the steel sheet as an observation surface, polished on the observation surface, and then etched into a nit, and from 1/8 to 3/3 centered on a 1/4 thickness from the surface of the plate thickness. 8 The range of thickness can be observed with an electric field emission scanning electron microscope (FE-SEM: Field Emission Scanning Electron Microscope) to measure the area fraction and take it as a volume fraction.

The volume fraction of retained austenite is evaluated by X-ray diffraction. In the range of 1/8 to 3/8 thickness from the surface of the plate thickness, the surface parallel to the plate surface is finished with a mirror surface, the area fraction of fcc iron is measured by X-ray diffraction method, and the retained austenite is taken with it. It can be considered as the volume fraction of.

The amount of solid solution carbon (Cγ [mass%]) in the retained austenite was subjected to an X-ray diffraction test under the same conditions as the measurement of the volume fraction of the retained austenite to obtain the average lattice constant a [nm] of the retained austenite, You can find it using an equation.

Cγ = 2.264 × 10 ² × (a-0.3556)

The aspect ratio of the retained austenite particles, the circle equivalent diameter, and the interface are evaluated by performing high resolution crystal orientation analysis by a transmission EBSD method (electron beam backscattering diffraction method) using FE-SEM. In the range of 1 / 8th to 3 / 8th thickness of the plate thickness, the thin flakes parallel to the plate surface are cut and subjected to mechanical polishing and electropolishing to the flakes to observe the periphery of the holes formed in the flakes, thereby allowing fine residual austenity. The night particles can be observed with high precision. In addition, "OIM Analysys 6.0" manufactured by TSL Corporation can be used for the analysis of the data obtained by the transmission EBSD method. Observation by the transmission EBSD method is performed by setting five or more areas of 2.0 × 10 ^-10 m 2 or larger in the flakes. From the observation results, the region judged to be fcc iron is referred to as residual austenite.

A method for obtaining the aspect ratio and the equivalent diameter of the retained austenite particles will be described. First, only fcc iron is extracted from the measured crystal orientation, and a crystal orientation map is drawn. A boundary that generates a crystal orientation difference of 10 ° or more is considered to be a grain boundary. The aspect ratio is the value obtained by dividing the long axis length of the particles by the short axis length. The circle equivalent diameter is obtained by obtaining an area from the number of measurement points included in each particle, multiplying the area by 4 / π, and then obtaining a square root.

A method for determining the ratio of the interface between the tempered martensite or fresh martensite occupied by the grain boundary of the retained austenite particles will be described. First, with respect to the data obtained from the bcc iron, "Grain Average Fit", which is an index indicating deformation of the crystal lattice in each crystal grain, is calculated using a boundary that generates a crystal orientation difference of 4 ° or more as a grain boundary.

The particles having a high grain average fit are tempered martensite or fresh martensite. For the sum StM + SfM [%] of the area fraction StM [%] of tempered martensite and the area fraction SfM [%] of fresh martensite obtained by observation using FE-SEM, the grain average fit is the horizontal axis, and the vertical axis To produce a histogram with the area of each crystal grain. In addition, the part corresponding to StM + SfM [%] from the side with high grain average fit is regarded as tempered martensite or fresh martensite.

Residual austenite, tempered martene, by setting the boundary which generates a crystal orientation difference of 4 ° on the crystallographic map of residual austenite as a grain boundary, and also marking an area considered as tempered martensite or fresh martensite on the same map The placement of the site and fresh martensite can be seen. Then, from this arrangement, the proportion occupied by the interface between the tempered martensitic or fresh martensite can be obtained from the total grain boundaries of the retained austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more.

Although the thickness of the high-strength steel sheet of the present embodiment is not limited, it is preferably 0.4 mm or more and 5.0 mm or less. When the plate thickness is less than 0.4 mm, it may be difficult to keep the shape of the steel sheet flat. When the plate thickness is more than 5.0 mm, it is difficult to control the heating conditions and cooling conditions in the manufacturing process, so that a predetermined microstructure may not be obtained.

(Second embodiment)

First, the high-strength galvanized steel sheet according to the second embodiment of the present invention will be described. The high-strength galvanized steel sheet according to the present embodiment has the high-strength steel sheet according to the first embodiment and a galvanized layer formed on its surface. The galvanized layer is, for example, a hot dip galvanized layer or an alloyed galvanized layer. The galvanized layer may be an electro-galvanized layer.

When a hot dip galvanizing layer is used, in order to improve the adhesion between the surface of the high-strength steel sheet and the hot dip galvanizing layer, the Fe content in the hot dip galvanizing layer is preferably 3.0% by mass or less. Further, the higher the Al content in the hot-dip galvanized layer, the more easily the adhesiveness between the surface of the high-strength steel sheet and the hot-dip galvanized layer tends to deteriorate, and the Al content is more than 0.5% by mass, so that the deterioration of adhesiveness is remarkable. Therefore, the Al content in the hot-dip galvanized layer is preferably 0.5% by mass or less.

When an alloyed galvanized layer is used, in order to increase the adhesion between the surface of the high-strength steel sheet and the alloyed galvanized layer, the Fe content in the alloyed galvanized layer is preferably 7.0% by mass or more and 13.0% by mass or less.

Galvanized layer, Ag, B, Be, Bi, Ca, Cd, Co, Cr, Cs, Cu, Ge, Hf, I, K, La, Li, Mg, Mn, Mo, Na, Nb, Ni, Pb, It may contain one or two or more selected from the group consisting of Rb, Sb, Si, Sn, Sr, Ta, Ti, V, W, Zr and REM. Depending on the content of each element, the corrosion resistance and workability of the galvanized layer may be improved.

Fe content and Al content in a galvanized layer can be measured by the method shown below. First, the plated layer is dissolved by immersing a plated steel sheet of a predetermined size in hydrochloric acid at a concentration of 3 to 10%, to which an inhibitor is added, at room temperature. Subsequently, the solution in which the plating layer is dissolved is diluted, and the concentrations of Zn, Al, and Fe in the diluted solution are analyzed by an inductively coupled plasma (ICP) method to obtain a mass ratio of Zn, Al, and Fe in the plating layer. Then, the Fe content and Al content in the plating layer are calculated from the mass ratio of Zn, Al and Fe in the plating layer. As an inhibitor, hexamethylenetetramine defined in JIS K 8847 may be used, for example.

Next, a method of manufacturing the high strength steel sheet according to the first embodiment will be described. This manufacturing method includes a hot rolling process, a pickling process, a cold rolling process, an annealing process, a bainite transformation process, a martensite transformation process, and a tempering process, and may include a softening annealing process. This manufacturing method is preferably applied to the production of high-strength steel sheets having a plate thickness of 0.4 mm or more and 5.0 mm or less. When the plate thickness is less than 0.4 mm, it may be difficult to keep the shape of the steel sheet flat. When the plate thickness is more than 5.0 mm, it is difficult to control the heating conditions and cooling conditions in the manufacturing process, so that a predetermined microstructure may not be obtained.

In order to manufacture the high-strength steel sheet of the first embodiment, first, a slab having the above-described chemical composition is cast. As the slab to be used for hot rolling, one produced by a continuous casting slab, a thin slab caster, or the like can be used. The slab after casting may be cooled to room temperature once, but it is more preferable to provide it to the hot rolling directly as it is at a high temperature, because energy required for heating can be reduced.

(Hot rolling process)

In the hot rolling process, if the heating temperature of the slab is less than 1080 ° C, coarse inclusions due to casting remain, and there is a possibility that the steel sheet breaks after hot rolling. For this reason, the heating temperature of the slab is 1080 ° C or higher, preferably 1150 ° C or higher. Although the upper limit of the heating temperature of the slab is not particularly determined, a large amount of energy is required to heat it in excess of 1300 ° C, so it is preferably 1300 ° C or less.

When the completion temperature of hot rolling is less than 850 ° C, the rolling reaction force becomes high, and it becomes difficult to stably obtain a desired plate thickness. For this reason, the completion temperature of hot rolling is set to 850 ° C or higher. From the viewpoint of rolling reaction force, the completion temperature of hot rolling is preferably 870 ° C or higher. On the other hand, in order to make the completion temperature of the hot rolling exceed 1020 ° C, an apparatus for heating the steel sheet is required from the end of heating of the slab to the completion of the hot rolling, and a high cost is required. For this reason, the completion temperature of hot rolling is set to 1020 ° C or lower. In addition, if the completion temperature of the hot rolling is excessively high, the shape of the steel sheet may collapse in the subsequent cooling process, and shape correction processing may be required after completion of the cooling, which is not preferable in view of cost. For this reason, the completion temperature of hot rolling is preferably 1000 ° C or lower, and more preferably 980 ° C or lower.

When the average cooling rate in the range from 850 ° C to 700 ° C is less than 8.0 ° C / sec between completion of hot rolling and winding, coarse ferrite is generated, Mn segregation becomes strong, and the microstructure with localized hard tissue Organization. As a result, the possibility that residual austenite in contact with tempering martensitic or fresh martensitic in the high-strength steel sheet is excessively increased. For this reason, in this method, the average cooling rate in the range from 850 ° C to 700 ° C is set to 8.0 ° C / sec or more. The average cooling rate in the range from 850 ° C to 700 ° C is preferably 12.0 ° C / sec or more in order to avoid the abutment of the retained austenite and the tempered martensite in the high-strength steel sheet to increase the impact resistance properties of the steel sheet. And more preferably 16.0 ° C / sec or more. Even if the upper limit of the average cooling rate in the range from 850 ° C to 700 ° C is not limited, the high-strength steel sheet according to the first embodiment can be produced.

When the temperature (winding temperature) T _C of winding the hot-rolled steel sheet as a coil is more than 700 ° C., phase transformation proceeds at high temperature during slow cooling to room temperature after coiling, coarse ferrite is generated, Mn segregation becomes strong, and hard tissue is formed. It becomes a localized microstructure. As a result, the possibility that residual austenite in contact with tempering martensitic or fresh martensitic in the high-strength steel sheet is excessively increased. For this reason, in this manufacturing method, the coiling temperature T _{C is set} to 700 ° C or lower. The coiling temperature T _C is preferably 660 ° C. or less in order to avoid the abutment of the retained austenite and tempering martensite in the high-strength steel sheet and to improve the impact resistance properties of the steel sheet. On the other hand, when the coiling temperature T _C is less than 400 ° C., the strength of the hot rolled steel sheet becomes excessively high, and there is a fear that the steel sheet breaks in the pickling and cold rolling processes. For this reason, the coiling temperature T _C is 400 ° C or higher. In order to cool the steel sheet with high precision, the coiling temperature T _C is preferably 500 ° C or higher.

If the cooling rate to the room temperature after winding is fast, the phase transformation after winding proceeds at a lower temperature, resulting in a microstructure in which the hard tissue is properly dispersed, and the retained austenite in the high-strength steel sheet is tempered martensite or fresh marten. The possibility of contact with the site is reduced. For this reason, the average cooling rate from the coiling temperature T _C to 350 ° _C is set to 5.0 x 10 ^-3 ° C / sec or more. On the other hand, if the steel sheet wound as a coil is cooled excessively and rapidly, a large temperature deviation occurs in the coil, and the phase transformation behavior also occurs with the temperature deviation. The possibility of residual austenite in contact with tempering martensite or fresh martensite among steel sheets increases. In order to avoid this, the average cooling rate from the coiling temperature T _C to 350 ° C. is 1.0 × 10 ⁻² ° C./sec or less, preferably 9.0 × 10 ⁻³ ° C./sec or less.

(Pickling process)

In the pickling process after the hot rolling process, pickling of the hot rolled steel sheet obtained in the hot rolling process is performed. In pickling, oxides present on the surface of the hot rolled steel sheet are removed. Pickling is important for improving the chemical conversion treatment and plating properties of the hot rolled steel sheet. Pickling of the hot rolled steel sheet may be performed only once, or may be performed multiple times.

(Cold rolling process)

In the cold rolling step after the pickling step, cold rolling of the hot rolled steel sheet is performed to obtain a cold rolled steel sheet. When the reduction ratio of the total in cold rolling is more than 85%, the ductility of the steel sheet is remarkably lowered, and there is an increased risk of breaking the steel sheet during cold rolling. For this reason, in cold rolling, the total reduction ratio is set to 85% or less, preferably 75% or less, and more preferably 70% or less. The lower limit of the reduction ratio of the total in the cold rolling step is not particularly determined, and cold rolling may be performed. However, when the total reduction ratio in cold rolling is less than 0.05%, the shape of the steel sheet becomes heterogeneous, the plating does not adhere uniformly, and the appearance may be damaged. For this reason, the rolling reduction ratio of the total in cold rolling is preferably 0.05% or more, and more preferably 0.10% or more. In addition, it is preferable to perform cold rolling in a plurality of passes, but the number of passes of cold rolling and the distribution of the reduction ratio for each pass are not asked.

If the reduction ratio of the total in cold rolling is more than 10% and less than 20%, recrystallization does not proceed sufficiently in the annealing process after the cold rolling process, and coarse grains that have lost their malleability, including a large amount of dislocation, remain near the surface layer. In some cases, the bendability and fatigue resistance may deteriorate. In order to avoid this, it is effective to reduce the total reduction ratio, to slightly reduce the accumulation of dislocations to the crystal grains, and to leave the malleability of the crystal grains effective. From the viewpoint of slightly reducing the accumulation of dislocations on the crystal grains, the reduction ratio of the total in the cold rolling step is preferably 10% or less, and more preferably 5.0% or less. In addition, in order to obtain good bending properties and fatigue resistance, the reduction ratio of the total in the cold rolling step is increased to sufficiently recrystallize in the annealing step, and the processed structure is made into a recrystallized grain with little accumulation of dislocation inside. It is valid. In order to sufficiently advance the recrystallization in the annealing step, the reduction ratio of the total is preferably 20% or more, and more preferably 30% or more.

(Softening annealing process)

The present production method may include a softening annealing process between the hot rolling process and the cold rolling process. In the soft nitridation annealing step, soft annealing for the purpose of softening the steel sheet is performed. By performing soft annealing, the rolling reaction force in cold rolling can be reduced, and the shape of the steel sheet can be improved. However, if the annealing temperature in the softening annealing process exceeds 680 ° C, the grain size of cementite becomes coarse, the frequency of nucleation of austenite particles in the annealing process decreases, and some austenite particles grow excessively. , The density of coarse austenite particles increases, and the impact characteristics deteriorate. For this reason, the annealing treatment temperature is preferably 680 ° C or lower. Although the lower limit of the annealing treatment temperature is not particularly determined, a sufficient softening effect cannot be obtained if it is less than 300 ° C. For this reason, the annealing treatment temperature is preferably 300 ° C or higher.

The soft nitridation annealing process may be either before or after the pickling process. Pickling may be performed both before and after the soft nitridation annealing step.

(Annealing process)

In the annealing process after the cold rolling process, annealing of the cold rolled steel sheet is performed. In the annealing process, is heated to a cold-rolled steel sheet obtained in the cold rolling process, the A _C1 point to (A _C1 +25 points), maximum heating temperature and the average heating rate above 0.5 ℃ / sec in the range of ℃ T _max. If the average heating rate in the range of A _C1 point to (A _C1 point +25) ° C. is less than 0.5 ° C./sec, the austenite particles produced immediately above the point A _C1 grow excessively, and the hard tissue is localized. It becomes micro organization. As a result, the possibility that residual austenite in contact with tempering martensitic or fresh martensitic in the high-strength steel sheet is excessively increased. The average heating rate in the range of A _C1 point to (A _C1 point +25) ° C. is preferably 0.8 in order to avoid the abutment of the retained austenite and tempering martensite in the high-strength steel sheet and to increase the impact resistance properties of the steel sheet. It is set to ℃ / sec or more. The upper limit of the average heating rate in the range of A _C1 point to (A _C1 point +25) ° C is not particularly set, but if the heating rate is excessively high, cementite in the steel will remain dissolved, and there is a risk of deterioration in properties. . For this reason, the average heating rate in the range of A _C1 point to (A _C1 point +25) ° C is preferably 100 ° C / sec or less.

When the maximum heating temperature T _max is less than (A _C1 point +40) ° C., cementite in the steel remains, and there is a fear that characteristics may deteriorate. For this reason, the maximum heating temperature T _max is set to (A _C1 point +40) ° C or higher. In order to increase the hard tissue fraction and make it higher in strength, the maximum heating temperature T _max is preferably (A _C1 point +55) ° C or higher. On the other hand, when the maximum heating temperature T _max is more than 1000 ° C, the austenite diameter becomes coarse, and the properties of the high-strength steel sheet may deteriorate significantly. For this reason, the maximum heating temperature T _max is preferably 1000 ° C or less.

The points A _C1 and A _C3 of the steel sheet are the starting point and the completion point of the austenite reverse transformation, respectively. The points A _C1 and A _{C3 of the} steel sheet are obtained by cutting small pieces from the steel sheet after hot rolling, heating it to 1200 ° C at 10 ° C / sec, and measuring the volume expansion during that time.

In the annealing step, the residence time t _max in the temperature range from the maximum heating temperature T _max to (T _max -10) ° C greatly affects the properties of the high-strength steel sheet. If this residence time t _max is too long, Mn segregation progresses, and the hard tissue becomes a localized microstructure. As a result, the likelihood of residual austenite in contact with tempering martensitic or fresh martensitic in the high-strength steel sheet is excessively high. Therefore, in this production method, in order to improve the impact resistance, the parameter Q1 expressed by (Equation 2) relating to the residence time t _max is set to 2.0 or less.

( _WMn in (Equation 2) is the content of _Mn in mass%, and T _C is the temperature [° C] in which the hot rolled steel sheet is coiled in a hot rolling process. T is T _max or A _C3 , whichever is lower. Temperature)

The parameter Q1 reflects the progress of the Mn segregation after the annealing process, so that the larger the parameter Q1, the Mn segregation proceeds, and the impact resistance characteristics deteriorate. From the viewpoint of impact resistance, the parameter Q1 is 2.0 or less, preferably 1.5 or less, and more preferably 1.0 or less.

After staying in the temperature range and the residence time t _max , the steel sheet is cooled to 650 ° C (first cooling). The average cooling rate (first cooling rate) up to 650 ° C (first cooling stop temperature) may be arbitrarily changed depending on the required characteristics of the steel sheet. The first cooling rate is preferably 0.5 ° C / sec or more. When the first cooling rate is 0.5 ° C / sec or more, generation of a large amount of pearlite can be prevented.

After the first cooling, the steel sheet is cooled with an average cooling rate of 2.5 ° C / sec or higher from 650 ° C to a temperature of 500 ° C or lower (second cooling stop temperature) (second cooling). If the average cooling rate (second cooling rate) in the temperature range from 650 ° C to 500 ° C is less than 2.5 ° C / sec, a large amount of pearlite or coarse cementite occurs, resulting in impaired formability. In addition, the ferrite produced in this temperature range is soft and reduces strength. In order to increase the strength of the high-strength steel sheet, the second cooling rate is preferably 5.0 ° C./sec or more, and more preferably 9.0 ° C./sec or more. Although the upper limit of the second cooling rate is not limited, the high-strength steel sheet according to the first embodiment can be produced, but in order to obtain an average cooling rate of 200 ° C / sec or more, it is necessary to employ a special cooling method. Therefore, from the viewpoint of cost, the second cooling rate is preferably 200 ° C / sec or less.

(Bainite transformation process)

In the bainite transformation step after the annealing step, a treatment that promotes bainite transformation is performed in a range of 500 to 340 ° C. By advancing the bainite transformation, the nucleation sites present in the particles of unmodified austenite are consumed, and the martensite produced in the particles of austenite in the next martensite transformation process is reduced. As a result, the interface between the retained austenite and the tempered martensite decreases, and the impact resistance properties are improved. On the other hand, if the bainite transformation proceeds excessively, carbon is excessively concentrated in the unused austenite. As a result, the Ms point in the unmodified austenite decreases, and martensite cannot be obtained in the next martensite transformation step. In order to appropriately advance the bainite transformation and improve the impact resistance, in the bainite transformation step, the parameter Q2 represented by (Expression 3) is set to 0.10 or more and 3.00 or less.

(In (Formula 3), Mn, Si, Cr, Al are the content of each element in mass%, t _B is the processing time [second] of the bainite transformation process, and Ti is from the start to completion of the bainite transformation process. It is the average temperature [° C] in the i-th range in which the treatment time is divided into 10 equal parts.)

The parameter Q2 reflects the progress of the bainite transformation, and the larger the parameter Q2, the progress of the bainite transformation. If the parameter Q2 is in the range of 0.10 or more and 3.00 or less, the progress of bainite transformation is appropriate. From the viewpoint of impact resistance, the parameter Q2 is preferably 0.25 or more and 2.50 or less.

In this manufacturing method, you may perform the process of cooling a steel plate to Ms *-(Mf * + 50) degreeC, and reheating it to 340-500 degreeC between an annealing process and a bainite transformation process. By this treatment, the nucleation site in the particles of unmodified austenite can be consumed more efficiently, and the impact resistance properties of the steel sheet can be further improved. Ms * is represented by (Formula 4), and Mf * is represented by (Formula 5).

(C, Mn, Ni, Cr, Si, Mo, Al in (Formula 4) is the content of each element in mass%, and is 0 when the element is not included. In (Formula 4) and (Formula 5) V _α is the volume fraction of ferrite [%])

In addition, it is difficult to directly measure the volume fraction of ferrite during the production of the steel sheet. For this reason, in this manufacturing method, a small piece of cold rolled steel sheet before an annealing process is cut out, and the small piece is annealed at the same temperature history as the annealing process to measure the change in ferrite volume of the small piece, and the numerical value calculated using the result , Used as the volume fraction of ferrite. When measuring the volume fraction of ferrite, when the steel sheet is manufactured under the same manufacturing conditions (temperature history), the results of the first measurement may be used, and it is not necessary to measure each time. When the manufacturing conditions are significantly changed, measurement is performed again. Of course, you may observe the microstructure of the steel plate actually manufactured and feedback to the manufacturing after the next time.

(Martensitic transformation process)

In the martensite transformation step after the bainite transformation step, a treatment for generating a martensite transformation is performed by cooling the steel sheet (third cooling). If the average cooling rate (third cooling rate) in the range of 340 ° C to Ms point during the third cooling is less than 1.0 ° C / sec, lower bainite containing carbides is generated, and formability deteriorates. Therefore, the third cooling rate is 1.0 ° C / sec or more. In order to obtain better moldability, the third cooling rate is preferably 2.5 ° C./sec or more, and more preferably 4.0 ° C./sec or more. The Ms point is represented by (Equation 6).

(In (Equation 6), C, Mn, Ni, Cr, Si, Al are the content of each element in mass%, and when the element is not included, it is 0. V _α is the volume fraction of ferrite [%], V _B is the volume fraction of bainite [%])

Then, it cools to arbitrary temperature (3rd cooling stop temperature) between Ms point or less and Mf point or more, and martensitic transformation is advanced. The Mf point is expressed by (Equation 7).

(In (Formula 7), C, Mn, Ni, Cr, Si, Al is the content of each element in mass%, and when the element is not included, it is 0. V _α is the volume fraction of ferrite [%], V _B is the volume fraction of bainite [%])

In addition, it is difficult to directly measure the volume fraction of ferrite and the volume fraction of bainite during the production of the steel sheet. For this reason, in this manufacturing method, a small piece of cold-rolled steel sheet before an annealing process is cut out, and the small piece is annealed at the same temperature history as the annealing process to measure the volume change of ferrite and bainite in the small piece, and using the result. The calculated values are used as the volume fraction of ferrite and the volume fraction of bainite, respectively. For the measurement of the volume fraction of ferrite and bainite, when manufacturing a steel sheet under the same manufacturing conditions (temperature history), the results of the first single measurement may be used, and it is not necessary to measure each time. When the manufacturing conditions are significantly changed, measurement is performed again. Of course, you may observe the microstructure of the steel plate actually manufactured and feedback to the manufacturing after the next time.

The third cooling stop temperature is preferably an Mf point of -10 ° C or less, and a range of -10 to 50 ° C. By setting the cooling stop temperature in the range of -10 to 50 ° C, it is possible to shift to the next tempering step without using a special heat keeping device or refrigerating device. Therefore, from the viewpoint of cost, the third cooling stop temperature is preferably in the range of -10 to 50 ° C.

In this manufacturing method, you may perform a 2nd cold rolling (skin pass rolling) between a martensitic transformation process and a tempering process. In the second cold rolling, cold rolling is performed on the steel sheet with a rolling rate of 3.0% or less, preferably 2.0% or less. By performing the second cold rolling, the unstable untransformed austenite particles having a large ratio of the interface formed with martensite occupying the entire grain boundary are transformed into martensite, thereby improving the impact resistance properties of the steel sheet after the tempering process.

(Tempering process)

In the tempering step after the martensitic transformation step, the steel sheet is tempered in the range of 200 to 600 ° C. By the tempering step, the martensite produced in the martensitic transformation step becomes a tempered martensite, and the moldability and impact resistance properties of the steel sheet are greatly improved.

Regarding the tempering temperature T _tem [° C] and the tempering treatment time t _tem [second] between the tempering temperature T _tem and (T _tem -10 ° C), the parameter Q3 represented by (Expression 8) is 1.0. The parameter Q4 represented by (Formula 9) is set to 1.00 or less, and the parameter Q5 represented by (Formula 10) is set to 1.00 or less.

(Si in (Equation 8), Si is the content of Si in mass%. T _tem is the tempering temperature [° C], and t _tem is the tempering treatment time [seconds].

The parameter Q3 represents the degree of tempering of martensite. If the parameter Q3 is less than 1.0, the tempering of martensite does not proceed sufficiently, and the hole expandability and stretch flangeability deteriorate. For this reason, the parameter Q3 needs to be 1.0 or more. In order to obtain better hole expandability and stretch flangeability, the parameter Q3 is preferably 1.5 or more, and more preferably 2.0 or more.

(In (Equation 9), Mn and Ni are the contents of Mn and Ni in mass%, and 0 is not included in the element. T _tem is the tempering temperature [° C], and t _tem is the tempering treatment time [seconds]. to be)

The parameter Q4 indicates the stability of residual austenite. When the parameter Q4 is more than 1.00, residual austenite is excessively stabilized, so that it does not transform into martensite during deformation, and the balance between strength and ductility deteriorates. For this reason, the parameter Q4 needs to be 1.00 or less. In order to further suppress the stabilization of the retained austenite, and to obtain a balance between superior strength and ductility, the parameter Q4 is preferably 0.90 or less, and more preferably 0.80 or less.

(In (Equation 10), Si, Al, Mn, Cr, Mo, and Ni are the content of each element in mass%, and when the element is not included, it is 0. T _tem is the tempering temperature [℃], t _tem Is the tempering treatment time [seconds])

The parameter Q5 represents the production behavior of pearlite and / or coarse cementite from residual austenite. When the parameter Q5 is more than 1.00, most of the retained austenite decomposes to pearlite and / or coarse cementite, deteriorating the strength and formability of the steel sheet. For this reason, the parameter Q5 needs to be 1.00 or less. In order to further suppress the decomposition of retained austenite, and to obtain superior strength and formability, the parameter Q5 is preferably 0.60 or less, and more preferably 0.20 or less.

In the tempering step, the average heating rate from 200 ° C to the highest heating temperature (tempering temperature) is preferably 1.0 ° C / sec or more. When the average heating rate is 1.0 ° C / sec or more, carbides in the tempered martensite become fine, and impact resistance characteristics can be improved. In order to improve the impact resistance properties, preferably, the average heating rate in the temperature range is more preferably 4.0 ° C./sec or more.

Further, after the tempering step, cold rolling with a rolling rate of 3.00% or less may be performed for the purpose of shape correction.

The high strength steel plate of 1st Embodiment is obtained by the above process.

Next, a method of manufacturing the high-strength galvanized steel sheet according to the second embodiment will be described. In this manufacturing method, the high-strength steel sheet according to the first embodiment obtained in the above-described manufacturing method or the high-strength steel sheet produced during the above-mentioned manufacturing method is subjected to plating treatment to form a galvanized layer. As the plating treatment, hot dip galvanizing treatment or electro galvanizing treatment may be performed.

Examples of the step (timing) of hot-dip galvanizing treatment on a high-strength steel sheet during production include timing before the bainite transformation process, timing before the martensite transformation process, timing before the tempering process, and after the tempering process. Further, during the tempering step, a hot dip galvanizing treatment may be performed.

The temperature of the zinc plating bath in the hot dip galvanizing treatment is preferably 450 to 470 ° C. When the temperature of the zinc plating bath is less than 450 ° C, the viscosity of the plating bath becomes excessively high, it becomes difficult to control the thickness of the plating layer, and the appearance of the steel sheet may be impaired. For this reason, the temperature of the zinc plating bath is preferably 450 ° C or higher. On the other hand, when the temperature of the galvanizing bath exceeds 470 ° C, a large amount of fume is generated, which makes it difficult to safely manufacture. For this reason, the temperature of the zinc plating bath is preferably 470 ° C or lower.

The temperature of the steel sheet immersed in the galvanizing bath (steel plate penetration temperature) is preferably 420 to 500 ° C. If the temperature of the steel sheet immersed in the galvanizing bath is less than 420 ° C, in order to stabilize the temperature of the galvanizing bath to 450 ° C or higher, it is necessary to apply a large amount of heat to the galvanizing bath, which is unsuitable for practical use. Therefore, in order to stabilize the bath temperature of the galvanizing bath, the steel sheet penetration temperature is preferably 420 ° C or higher, and more preferably 440 ° C or higher. On the other hand, if the temperature of the steel plate immersed in the galvanizing bath exceeds 500 ° C, in order to stabilize the temperature of the galvanizing bath to 470 ° C or lower, it is necessary to introduce a facility for heat sinking a large amount of heat from the galvanizing bath. , In terms of manufacturing cost. Therefore, in order to stabilize the bath temperature of the galvanizing bath, the steel sheet intrusion temperature is preferably 500 ° C or lower, and more preferably 480 ° C or lower.

The zinc plating bath preferably has a composition in which zinc is mainly used and the effective Al content, which is a value obtained by subtracting the total Fe content from the total Al content in the zinc plating bath, is 0.010 to 0.300 mass%. If the effective Al content in the galvanizing bath is less than 0.010% by mass, the penetration of Fe into the galvanizing layer proceeds excessively, and the adhesion to plating may be impaired. Therefore, the effective Al content in the galvanizing bath is preferably 0.010 mass% or more, more preferably 0.030 mass% or more, and even more preferably 0.050 mass% or more. On the other hand, if the effective Al content in the galvanizing bath is more than 0.300% by mass, Fe-Al intermetallic compound is excessively formed at the boundary between the ferrous iron and the galvanizing layer, and plating adhesion may be remarkably impaired. Therefore, the effective Al content in the zinc plating bath is preferably 0.300 mass% or less. In particular, when the alloying treatment is performed, when Fe-Al intermetallic compounds are formed, the movement of Fe and Zn atoms is inhibited, and the formation of the alloy phase is suppressed, so that the effective Al content in the plating bath is more preferably 0.180 mass% or less. And more preferably 0.150% by mass or less.

The galvanized bath is Ag, B, Be, Bi, Ca, Cd, Co, Cr, Cs, Cu, Ge, Hf, I, K, La, Li, Mg, Mn, Mo, Na, Nb, Ni, Pb , Rb, Sb, Si, Sn, Sr, Ta, Ti, V, W, Zr, and REM. Depending on the content of each element, the corrosion resistance and workability of the galvanized layer may be improved.

After immersing the steel sheet in a galvanizing bath, in order to obtain an appropriate plating adhesion amount, preferably, a high pressure gas mainly containing nitrogen is sprayed on the surface of the steel sheet to remove excess plating solution on the surface layer.

Examples of the step (timing) in which the high-strength steel sheet during manufacturing is subjected to an electro-galvanization treatment include timing before the tempering process and after the tempering process. As the electro-galvanizing treatment, a conventionally known method can be used. As the electro-galvanizing bath, for example, it may be mentioned H ₂ SO _4, ZnSO ₄ and NaSO ₄ pH 1.5 to 2.0 comprising a. Conditions such as temperature and current density of the electro-galvanizing bath can be appropriately determined depending on the type of the electro-galvanizing bath, the thickness of the galvanized layer, and the like. In addition, when the high-strength steel sheet is subjected to an electro-galvanization treatment, preferably, the high-strength steel sheet is pickled and then immersed in an electro-galvanization bath. As the pickling method of the high-strength steel sheet, a known method can be used, and for example, a method of immersing the high-strength steel sheet in sulfuric acid and performing bubbles until hydrogen bubbles are visually confirmed.

When an alloyed galvanized layer is formed as a galvanized layer, an alloying treatment is performed after plating treatment to form an alloyed zinc plated layer. The alloying treatment may be performed at any stage (timing) as long as it is after the plating treatment, or may be continuously performed after the plating treatment. Specifically, for example, when the plating treatment is performed before the bainite transformation step, as a step (timing) of alloying treatment, the timing before the bainite transformation step after the plating treatment, the timing before the martensite transformation step, the timing before the tempering step, Simultaneous with the tempering treatment, after a tempering process, etc. are mentioned. For example, in the case where the plating treatment is performed before the martensitic transformation process, the step of performing an alloying treatment (timing) includes the timing before the martensitic transformation process after the plating treatment, the timing before the tempering process, and the same time as the tempering process, after the tempering process. Can be lifted. For example, when the plating treatment is performed before the tempering step, the step (timing) of the alloying treatment includes timing before the tempering step after the plating treatment, simultaneous with the tempering treatment, and after the tempering step.

When the alloying treatment is performed after the hot dip galvanizing treatment is performed, for example, a heat treatment is preferably performed at a temperature of 470 to 600 ° C for 2 to 100 seconds. When the alloying treatment is performed after the electro-galvanizing treatment, for example, it is preferable to perform a heat treatment, which is preferably held at a temperature of 400 to 600 ° C for 2 to 100 seconds.

The obtained high-strength galvanized steel sheet may be cold rolled at a rolling reduction of 3.00% or less for shape correction.

Through the above steps, the high-strength galvanized steel sheet of the second embodiment is obtained.

The steel sheet having the chemical composition of the above-described high-strength steel sheet and manufactured through the treatment up to the tempering step in the same manner as the above-described method of manufacturing the high-strength steel sheet can be used as a steel sheet for tempering treatment. The steel sheet for tempering treatment may have a zinc plating layer on the surface. The galvanized layer can be formed by performing a plating treatment in the same way as the above-described method for producing a high-strength galvanized steel sheet before the tempering step.

The microstructure of the steel sheet for tempering treatment is as follows, for example.

(Micro organization)

The steel sheet for tempering treatment is in the range of 1/8 thickness to 3/8 thickness centering on 1/4 thickness from the surface of the plate thickness, in a volume fraction, ferrite: 85% or less, bainite: 3% or more 95% or less, fresh martensite: 1% or more and 80% or less, retained austenite: 1% or more and 25% or less, and also has microstructures represented by pearlite and cementite: 5% or less in total.

The reason for limiting each volume fraction of the structures other than fresh martensite is the same as that of the high-strength steel sheet according to the first embodiment. The reason for limiting the volume fraction of fresh martensite is the same as the volume fraction of tempered martensite in the high-strength steel sheet according to the first embodiment. This is because almost all of the fresh martensite becomes tempered martensite by a tempering process.

In the steel sheet for tempering treatment, the solid carbon content in the retained austenite is 0.60 to 0.95% by mass. When the solid carbon content in the retained austenite is within this range, a high-strength steel sheet having a solid carbon content in the retained austenite of 0.70 to 1.30 mass% is obtained through a tempering process.

In addition, this invention is not limited to the said embodiment. For example, as the plating layer, a nickel plating layer may be formed instead of a zinc plating layer. A coating made of a phosphorus oxide and / or a composite oxide containing phosphorus may be formed on the galvanized layer. Such a coating can function as a lubricant when processing a steel sheet, and can protect the galvanized layer. A coating made of phosphorus oxide and / or a composite oxide containing phosphorus can be formed using a known method.

In addition, all of the above-described embodiments are merely examples of specificization in carrying out the present invention, and the technical scope of the present invention should not be limitedly interpreted by them. That is, the present invention can be implemented in various forms without departing from its technical spirit or its main features.

Example

Next, examples of the present invention will be described. The conditions in Examples are one conditional examples employed to confirm the feasibility and effect of the present invention, and the present invention is not limited to these one conditional examples. The present invention is capable of employing various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.

Slabs having chemical compositions A to AK shown in Tables 1 and 2 were cast, heated to the slab heating temperatures shown in Tables 3 and 4, and subjected to hot rolling, cooling, coiling, and coil cooling to obtain hot rolled steel sheets (hot rolled). fair). Slab heating temperature, rolling completion temperature of hot rolling, average cooling rate in the range of 850 to 700 ° C (average cooling rate before winding), average cooling rate from winding temperature T _C and winding temperature T _C to 350 ° _C ( The coil cooling rate) is shown in Table 3 and Table 4. The blanks in Table 1 and Table 2 indicate that the content of the element was below the detection limit, and the balance is Fe and impurities. The underline in Table 1 and Table 2 indicates that the values are out of range.

Thereafter, the hot rolled steel sheet was pickled and cold rolled to obtain a cold rolled steel sheet (cold rolling process). Table 3 and Table 4 show the total rolling reduction of cold rolling. Some of the hot rolled steel sheets were subjected to cold rolling after performing a soft annealing process at the processing temperatures shown in Tables 3 and 4. Underlined in Table 3 and Table 4 indicates that the value is outside the range required for the production of the high-strength steel sheet according to the present invention.

Next, annealing of the cold rolled steel sheet was performed (annealing step). Table 5 shows the average heating rate in the range of A _C1 to (A _C1 +25) ° C of annealing, the maximum heating temperature T _max , and the maximum heating temperature T _max to the residence time in the range (T _max -10) from Table 5 and It is shown in Table 6. Table 5 also shows the average cooling rate (first cooling rate) of the first cooling to 650 ° C after the residence time at the above residence time and the average cooling rate (second cooling rate) in the range of 650 ° C to 500 ° C for the second cooling. And Table 6. In Table 5 and Table 6, A _C1 point, A _C3 point, T _max -A _C1 point, and parameter Q1 are also shown. The underline in Table 5 and Table 6 indicates that the numerical value is outside the range required for the production of the high-strength steel sheet according to the present invention.

The 2nd cooling was stopped at the 2nd cooling stop temperature, and the bainite transformation process was performed (bainite transformation process). Table 2 and Table 8 show the second cooling stop temperature and the average treatment temperature T _B of the bainite transformation treatment, treatment time t _B , and parameter Q2. The average processing temperature T _B is calculated using the formula shown below.

T _B = Σ (T _i ) / 10 = (T ₁ + T ₂ + T ₃ +… + T ₁₀ ) / 10

(Wherein, T _i represents the average temperature [° C] in the i-th range obtained by dividing the treatment time of the bainite transformation process into 10 equal parts)

For the experimental example in which the second cooling stop temperature was in the range of Ms * to Mf * + 50 [° C] shown in Tables 7 and 8, after cooling was stopped at the second cooling stop temperature, reheating was performed to a temperature of 340 to 500 ° C. Then, bainite transformation was performed. Underlined in Table 7 and Table 8 indicates that the value is outside the range required for the production of the high-strength steel sheet according to the present invention.

After the bainite transformation treatment, a martensitic transformation treatment was performed (martensitic transformation process). Table 9 and Table 10 show the average cooling rate (third cooling rate) and stop temperature (third cooling stop temperature) in the range of 340 ° C to Ms point of the third cooling in the martensite transformation process. After the martensitic transformation process, some cold rolled steel sheets were subjected to second cold rolling at the rolling rates shown in Tables 9 and 10 (second cold rolling process). Ms point and Mf point are also shown in Table 9 and Table 10 together. Underlined in Table 9 and Table 10 indicates that the value is outside the range required for the production of the high-strength steel sheet according to the present invention.

In this way, various steel sheets for tempering treatment were produced. In addition, with respect to the steel plates for tempering treatment of Experimental Examples No. 1 to No. 101, "ferrite (α)" and "bainite (B) by the method similar to the method of measuring the volume fraction of each structure of the high-strength steel plate mentioned above. ) ”,“ Tempered martensite (tM) ”,“ Residual austenite (residual γ) ”,“ Fresh martensite (fM) ”,“ Perlite and cementite total (P + C) ”volume fraction of each tissue Investigated. The amount of carbon dissolved in the retained austenite was examined for the steel sheets for tempering treatment of Experimental Examples No. 1 to No. 101 by the above-described method. Table 11 and Table 12 show the results. Underlined in Table 11 and Table 12 indicates that the value is outside the range required for the production of the high-strength steel sheet according to the present invention.

After the martensitic transformation step, a tempering treatment was performed (tempering step). Table 13 and Table 14 show the tempering temperature T _tem in the tempering process, the treatment time t _tem and the average heating rate from 200 ° C. to the tempering temperature T _tem . Tables 13 and 14 also show parameters Q3 to Q5. Underlined in Table 13 and Table 14 indicates that the value is outside the range required for the production of the high-strength steel sheet according to the present invention.

Experimental Examples No.13, No.22, No.25, No.28, No.34, No.37, No.41, No.43, No.49, No.50, No.60, No.61, In No.67, No.70, No.73, No.77, No.81, and No.83, the hot-dip galvanizing treatment was performed on the steel sheet in the steps shown in Table 15 (timing of plating treatment). Table 15 shows the steel sheet penetration temperature of the hot dip galvanizing treatment, the plating bath temperature, and the effective Al amount of the plating bath. In Examples No. 13, No. 22, No. 28, No. 34, No. 41, No. 50, No. 60, No. 61, No. 70, No. 73 and No. 83, Table 15 The alloying treatment was performed in the steps shown (timing of alloying treatment). Table 15 shows the temperature and the holding time of the alloying treatment. In the column of "surface" in Table 15, "GI" means that a hot-dip galvanized layer was formed on the surface, and "GA" means that an alloyed zinc-plated layer was formed on the surface.

In Experimental Example No. 2, in the step after the tempering treatment (timing of the plating treatment), the steel sheet was subjected to the electro zinc plating treatment shown below. In Example No. 65, in the step after the tempering treatment (timing of the plating treatment), the steel sheet was subjected to an electro-galvanizing treatment as shown below, and in the step after the tempering treatment (timing of the alloying treatment), an alloying treatment was performed.

The electro-galvanization treatment in Experimental Example No. 2 and Experimental Example No. 65 was performed by a method of energizing by immersing the steel sheet in an electrogalvanizing bath after pickling. The pickling was immersed in 10% sulfuric acid at room temperature, and carried out until bubbles of hydrogen were visually observed. Electrogalvanization is performed in an electroplating bath having a pH of 1.5 to 2.0 and a liquid temperature of 50 ° C., including H ₂ SO ₄ of 1.5 g / L, ZnSO ₄ of 194 g / L, and NaSO ₄ of 45 g / L. Was 25A / (dm) ² .

In Experimental Examples No. 3 to No. 35 and No. 60 to No. 80, after tempering treatment, cold rolling with a maximum rolling reduction of 1.00% was performed.

In Experimental Examples No. 22 and No. 83, after the tempering treatment, a degreasing agent was applied and washed with water spray. Subsequently, a complex oxide containing phosphorus oxide and phosphorus on the surface by a method of immersing in the order of surfactant, chemical conversion treatment liquid PB-SX35 manufactured by Nippon Parkerizing Co., Ltd., followed by washing with water spray and drying in a hot air oven. A coating film consisting of was formed.

About the experimental example in which the test was not stopped among the high-strength steel sheets and high-strength galvanized steel sheets of Experimental Examples No. 1 to No. 101 thus obtained, "ferrite (α)" and "vanite (B)" ”,“ Tempered martensite (tM) ”,“ Residual austenite (residual γ) ”,“ Fresh martensite (fM) ”,“ Pullite and cementite total (P + C) ”volume fraction of each tissue were investigated. did. The results are shown in Table 16 and Table 17. The underline in Table 16 and Table 17 indicates that the value is out of range.

In addition, for the experimental examples in which the test was not stopped among the high-strength steel sheets of Experimental Examples No. 1 to No. 101 and the high-strength galvanized steel sheet, the amount of solid carbon in the retained austenite was examined by the above-described method. Further, for the experimental example in which the test was not stopped, by the above-described method, among the total grain boundaries of the retained austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more, the tempered martensite or fresh martensite The proportion occupied by the interface (the proportion of the interface) was examined. For the experimental example in which the test was not stopped, the density of the retained austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more was examined. The results are shown in Table 18 and Table 19. Underlined in Table 18 and Table 19 indicates that the value is out of the scope of the present invention.

About the high-strength galvanized steel sheet in which the hot-dip galvanized layer was formed in Experimental Examples No. 1 to No. 101, Fe content and Al content in the plated layer were investigated by the above-described method. The results are shown in Table 18 and Table 19. In addition, with respect to the galvanized steel sheet for tempering treatment in which the hot dip galvanized layer was formed before the tempering step, the Fe content and Al content in the plating layer were investigated in the same manner as the high-strength galvanized steel sheet. The results are shown in Table 2 0 and Table 21. In Tables 18 to 21, "GI" means that a hot dip galvanized layer was formed on the surface, "GA" means that an alloyed zinc plated layer was formed on the surface, and "EG" formed an electro galvanized layer on the surface. Means, and "CR" means that no plating layer was formed on the surface.

The yield strength YS, the maximum tensile strength TS, the hole expandability λ, the ductility El, and the Charpy impact properties were investigated for the high-strength steel sheets and high-strength galvanized steel sheets of Experimental Examples No. 1 to No. 101 by the methods shown below. . The results are shown in Tables 22 to 25. The underlining in Tables 22 to 25 indicates that the desired value is out of range.

In the tensile test, a test piece No. 5 described in JIS Z 2201 was produced from a steel sheet, and the method was described in JIS Z2241 to yield yield strength YS, tensile maximum strength TS, and total elongation El. The hole expansion test was performed by the method described in JIS Z 2256. The ductility (total elongation) El and the hole expandability λ change with the tensile maximum strength TS, but make the strength, ductility and hole expandability favorable when satisfying the following expressions (11) and (12).

TS × El≥1.5 × 10 ⁴ … (Equation 11)

(YS × TS) ^0.75 × El × λ ^0.5 ≥2.0 × 10 ⁶ … (Equation 12)

Charpy impact properties are obtained in the direction perpendicular to the rolling direction (C direction) by collecting a plate of the same size as the impact test piece specified in JIS Z 2242, stacking three sheets and fastening with bolts, and then in the method specified in JIS Z 2242. Charpy absorption energy at 20 ° C (vE ₂₀ ) and Charpy absorption energy at -20 ° C (vE _-20 ) were measured. When the following (Equation 13) is satisfied, and both vE ₂₀ and vE- ₂₀ are ₂₀ J / cm 2 or more, the impact resistance properties are improved.

(vE _-20 ) / (vE ₂₀ ) ≥0.65… (Equation 13)

In addition, the closer the value of the left side of (Formula 13) is to 1.0, the better the impact resistance property is, and it is preferable that the value of (Formula 13) is 0.75 or more.

Experimental example No. 91 is an example in which moldability was deteriorated because the content of C was small and residual austenite was not obtained.

In Experimental Example No. 92, since C contained excessively, the fraction of retained austenite became excessive, and the impact characteristics deteriorated.

Experimental Example No. 93 is an example in which residual austenite was not obtained because the contents of Si, Mn, and Al were small and the parameter Q0 was excessive, so that characteristics were deteriorated.

As an example in which Experimental Example No. 94 contained Si excessively, the test was stopped because the properties of the slab deteriorated and the slab broke in the casting process.

Experimental Example No. 95 is an example in which residual austenite was not obtained because the content of Mn was small and pearlite was generated due to insufficient quenchability, and the characteristics were deteriorated.

Experimental Example No. 96 was an example in which Mn was contained excessively, and the test was stopped because the characteristics of the slab deteriorated and the slab broke during heating in a hot rolling process.

Experimental Example No. 97 was an example in which P was contained excessively, and the characteristics of the slab deteriorated, and the test was stopped because the slab was broken during transportation from the casting process to the hot rolling process.

Experimental Example No. 98 is an example in which S is excessively contained. Since coarse sulfide is generated in a large amount in steel, the formability and impact resistance of the steel sheet are greatly deteriorated.

In Experimental Example No. 99, as an example in which Al was contained excessively, the characteristics of the slab deteriorated, and the test was stopped because the slab broke in the casting process.

Experimental Example No. 100 is an example in which N is excessively contained. Since coarse nitride is generated in a large amount in steel, the formability and impact resistance of the steel sheet are greatly deteriorated.

Experimental Example No. 101 is an example in which O is contained excessively, and since a large amount of coarse oxide is generated in steel, the formability and impact resistance of the steel sheet are greatly deteriorated.

In Experimental Example No. 84, the heating temperature of the slab in the hot-rolling process was low, and the moldability of the hot-rolled steel sheet was remarkably impaired, and the test was stopped because the steel sheet fractured in the cold rolling process.

In Experiment No. 54, the rolling completion temperature in the hot rolling process was low, and the shape of the hot rolled steel sheet was significantly damaged, so the test was stopped.

In Experimental Example No. 75, the average cooling rate in the range of 850 ° C to 700 ° C in the hot rolling process was low, the localization of the hard structure became strong, and the ratio of the interface with the tempered martensite of residual austenite increased, resulting in impact characteristics. This is an example of deterioration.

Experimental Example No. 57 is an example in which the cooling rate of the coil in the hot rolling process is low, the localization of the hard structure becomes strong, the ratio of the interface with the tempered martensite of residual austenite increases, and the impact characteristics deteriorate.

In Experimental Examples No. 103 and No. 105, the cooling rate of the coil in the hot rolling process is high, the localization of the hard structure becomes strong, and the ratio of the interface formed with the tempered martensite of the retained austenite increases, resulting in deterioration of impact characteristics. Yes.

Experimental Example No. 106 is an example in which the coiling temperature T _C in the hot rolling process is high, the Mn localization becomes strong, the ratio of the interface formed with the tempered martensite of residual austenite increases, and the impact characteristics deteriorate.

In Experimental Example No. 42, the average heating rate in the range of A _C1 point to (A _C1 point +25) ° C. in the annealing step was low, the localization of the hard tissue became strong, and the residual austenite and tempered martensite This is an example in which the impact ratio is deteriorated because the interface ratio is high and the residual austenite particles are coarsened.

In Experimental Example No. 87, the highest heating temperature T _max in the annealing process was low, and since the coarse carbide was dissolved in a large amount, the total ratio of pearlite and cementite was increased, resulting in deterioration of formability.

In Experimental Examples No. 21 and No. 69, the parameter Q1 related to the maximum heating temperature T _max and the residence time t _max in the annealing process was excessive, so that the localization of the hard tissue became strong, and the tempered martensite of the retained austenite was achieved. This is an example in which the proportion of the interface is high and the impact characteristics are deteriorated.

In Experimental Example No. 78, the average cooling rate (first cooling rate) in the range of (highest heating temperature -10 ° C) of the annealing process to 650 ° C was low, and a large amount of soft pearlite was generated during cooling, resulting in steel sheet This is an example in which the balance of strength and formability of the product deteriorates.

In Experimental Example No. 36, the average cooling rate (second cooling rate) in the range of 650 ° C. to 500 ° C. in the annealing process was low, and a large amount of soft pearlite was generated during cooling to balance the strength and formability of the steel sheet. Is an example of deterioration.

In Experimental Example No. 12, the parameter Q2 related to the average treatment temperature T _i and the treatment time T _B of the bainite transformation process was insufficient, so that bainite transformation did not proceed sufficiently, and tempering martensite and residual austenite were not obtained. It is an example in which moldability deteriorates.

In Experimental Example No. 45, the parameter Q2 of the average treatment temperature T _i and the treatment time T _B of the bainite transformation process was excessive, and an example in which the bainite transformation proceeded excessively, tempering martensite and residual austenite were not obtained , The balance of strength and formability deteriorates.

In Experimental Example No. 27, the average cooling rate from the bainite transformation process to the martensitic transformation process, that is, the third cooling rate was low, a large amount of lower bainite was generated, tempering martensite was not obtained, and fresh martensite There are many, and this is an example in which moldability deteriorates.

Experimental Example No. 6 is an example in which the cooling stop temperature in the martensitic transformation process was lower than the Mf point, and residual austenite was not obtained, and the moldability was deteriorated.

In Experimental Example No. 24, the cooling stop temperature in the martensitic transformation process was higher than the Ms point.Temperature martensite was not obtained, and part of the austenite was transformed into fresh martensite after tempering treatment. This is an example in which the balance of formability and impact resistance are deteriorated.

Experimental Example No. 53 is an example in which excessive cold rolling was performed between the martensitic transformation process and the tempering treatment process, and the residual austenite formed in the martensitic transformation process is transformed by rolling, and the residual austenite is lost after the tempering treatment. , It is an example in which the balance of strength and formability deteriorates.

In Experimental Example No. 15, since the tempering temperature T _tem in the tempering treatment process was less than 200 ° C, martensite was not sufficiently tempered, and fresh martensite remained in a large amount, so that the balance of strength and formability and impact resistance characteristics were obtained. This is an example of deterioration.

In Experimental Example No. 18, the parameter Q3 related to the tempering temperature T _tem and the processing time t _tem in the tempering treatment step was too small, and martensite was not sufficiently tempered, and fresh martensite remained in a large amount, thereby increasing strength and molding. This is an example of deterioration of the balance and impact resistance of the castle.

In Experimental Example No. 48, the parameter Q4 for the tempering temperature T _tem and the treatment time t _tem in the tempering treatment process was excessive, and the amount of solid solution C of the retained austenite increased excessively, and the balance between strength and formability deteriorated. to be.

In Experimental Examples No. 33 and No. 66, the parameter Q5 for the tempering temperature T _tem and the processing time t _tem in the tempering treatment process is excessive, and a large amount of pearlite is generated, resulting in deterioration in the balance between strength and formability. .

In Experimental Examples No. 15, No. 18, No. 33, No. 48, and No. 66, all desired steel plates for tempering were obtained, but the subsequent tempering treatment was inadequate, and the microstructure was not properly controlled. , This is an example where the characteristics after tempering become inferior.

Experimental Examples No. 1 to No. 5, No. 7 to No. 11, No. 13, No. 14, No. 16, No. 17, No. 19, No. 20, No. 22, No. 23, No.25, No.26, No.28 to No.32, No.34, No.35, No.37 to No.41, No.43, No.44, No.46, No.47, No. 49 to No.52, No.55, No.56, No.58 to No.62, No.64, No.65, No.67, No.68, No.70 to No.74, No.76, No.77, No.79 to No.83, No.85, No.86, No.88 to No.90, No.102, No.104 according to the present invention have high strength excellent in moldability and impact resistance This is an example in which a steel sheet is obtained.

Experimental Example No. 29 is an example in which a high-strength steel sheet excellent in moldability and impact resistance is obtained by performing a soft nitridation treatment between a hot rolling process and a cold rolling process.

Experimental Example No. 3 is an example in which a high-strength steel sheet excellent in moldability and impact resistance is obtained by performing soft nitridation treatment between a hot-rolling process and a cold-rolling process, and also performing cold rolling between a martensitic transformation process and a tempering process. .

In Experimental Examples No. 46 and No. 52, the cooling stop temperature before the bainite transformation process, that is, the second cooling stop temperature was set to a range of Ms * point or less and Mf * point or more, and cooling was stopped at the second cooling stop temperature. This is an example in which a bainite transformation is performed after reheating to a temperature of 340 to 500 ° C, and a part of the microstructure is transformed into martensite before bainite transformation. Experimental examples No. 46 and No. 52 are examples in which a high-strength steel sheet having excellent impact resistance properties is obtained.

In Experimental Examples No. 4, No. 9, No. 10, No. 17, No. 19, No. 30, and No. 79, cold rolling was performed between the martensitic transformation process and the tempering process to formability and impact resistance. This is an example of obtaining a high-strength steel sheet having excellent properties.

Experimental Example No. 2 is an example in which a high-strength galvanized steel sheet excellent in moldability and impact resistance is obtained by performing cold rolling between a martensitic transformation step and a tempering step and electroplating after the tempering treatment.

In Experimental Examples No.25, No.43, No.49, and No.67, between the annealing process and the bainite transformation process, the steel sheet was immersed in a galvanizing bath to provide a high-strength galvanized steel sheet with excellent moldability and impact resistance. This is an example obtained.

Experimental examples No.77 and No.81 are examples in which a high-strength galvanized steel sheet excellent in moldability and impact resistance is obtained by immersing the steel sheet in a galvanizing bath between the bainite transformation process and the martensite transformation process.

Experimental Example No. 37 is an example in which a high-strength galvanized steel sheet excellent in moldability and impact resistance is obtained by immersing the steel sheet in a galvanizing bath in a tempering treatment step.

Experimental Example No. 50 is an example in which a steel sheet was immersed in a galvanizing bath between an annealing process and a bainite transformation process, and immediately subjected to an alloying treatment to obtain a high-strength galvanized steel sheet excellent in moldability and impact resistance.

In Experimental Examples No. 22 and No. 83, the steel sheet was immersed in a galvanizing bath between an annealing process and a bainite transformation process, and immediately subjected to an alloying treatment, which contained phosphorus oxide and phosphorus on the plating layer after the tempering treatment. It is an example in which a high-strength galvanized steel sheet excellent in moldability and impact resistance is obtained by applying a film made of a composite oxide.

In Experimental Example No. 60, a softening treatment was performed between the hot-rolling step and the cold-rolling step, the steel plate was immersed in a galvanizing bath between the annealing step and the bainite transformation step, and then immediately subjected to alloying treatment to formability. And high-strength galvanized steel sheet excellent in impact resistance.

In Experimental Examples No.34, No.41, and No.73, the steel sheet was immersed in a galvanizing bath between a bainite transformation process and a martensitic transformation process, and immediately subjected to an alloying treatment, resulting in moldability and impact resistance properties. This is an example in which an excellent high-strength galvanized steel sheet is obtained.

Experimental Example No. 13 is an example in which a high-strength galvanized steel sheet excellent in moldability and impact resistance is obtained by immersing the steel sheet in a zinc plating bath in the tempering treatment step, and further performing an alloying treatment.

In Experimental Example No. 70, a steel sheet was immersed in a galvanizing bath between an annealing process and a bainite transformation process, and an alloying treatment was performed between the bainite transformation process and the martensite transformation process to formability and impact resistance. This is an example in which an excellent high-strength galvanized steel sheet is obtained.

In Experimental Example No. 61, the steel sheet was immersed in a galvanizing bath between the annealing process and the bainite transformation process, and the cooling stop temperature before the bainite transformation process, that is, the second cooling stop temperature was Ms * point or less and Mf * point It is an example in which a high-strength galvanized steel sheet having excellent impact resistance properties is obtained by performing alloying treatment between the bainite transformation step and the martensite transformation treatment step in the above range.

In Experimental Example No. 28, the steel sheet was immersed in a galvanizing bath between the bainite transformation process and the martensitic transformation process, and subjected to an alloying treatment in the tempering treatment process, resulting in high strength zinc plating excellent in moldability and impact resistance. This is an example in which a steel sheet is obtained.

In Experimental Example No. 65, a high-strength galvanized steel sheet excellent in moldability and impact resistance was obtained by subjecting the steel sheet to an electroplating process between the martensitic transformation process and the tempering process, and performing an alloying process in the tempering process. This is an example obtained.

The present invention can be used, for example, in industries related to high-strength steel sheets such as automobile steel sheets and high-strength galvanized steel sheets.

Claims (8)

In mass%,
C: 0.075 to 0.400%,
Si: 0.01 to 2.50%,
Mn: 0.50 to 3.50%,
P: 0.1000% or less,
S: 0.0100% or less,
Al: 2.000% or less,
N: 0.0100% or less,
O: 0.0100% or less,
Ti: 0.000 to 0.200%,
Nb: 0.000 to 0.100%,
V: 0.000 to 0.500%,
Cr: 0.00 to 2.00%,
Ni: 0.00 to 2.00%,
Cu: 0.00 to 2.00%,
Mo: 0.00 to 1.00%,
B: 0.0000 to 0.0100%,
W: 0.00 to 2.00%,
Ca, Ce, Mg, Zr, La and one or more selected from the group consisting of REM: 0.0000 to 0.0100% in total,
Balance: Fe and impurities, also
Parameter Q0 represented by (Equation 1): 0.35 or more
Has a chemical composition represented by,
In the range of 1/8 thickness to 3/8 thickness centering on 1/4 thickness from the surface of the plate thickness, in a volume fraction,
Ferrite: 85% or less,
Bainite: 3% or more and 95% or less,
Tempering martensite: 1% or more and 80% or less,
Residual austenite: 1% or more and 25% or less,
Pearlite and coarse cementite: 5% or less in total, and
Fresh martensite: 5% or less
Has a micro-organism represented by,
The dissolved carbon content in the retained austenite is 0.70 to 1.30% by mass,
A high-strength steel sheet having an aspect ratio of 2.50 or less, and a ratio of the interface between the tempered martensite or fresh martensite and 75% or less of the total grain boundaries of the retained austenite particles having a circle equivalent diameter of 0.80 µm or more.
Q0 = Si + 0.1Mn + 0.6Al ... (Equation 1)
(In (Equation 1), Si, Mn and Al are the content of each element in mass%) According to claim 1,
In mass%,
Ti: 0.001 to 0.200%,
Nb: 0.001 to 0.100% and
V: 0.001 to 0.500%
High-strength steel sheet, characterized in that it contains one or two or more selected from the group consisting of. According to claim 1,
In mass%,
Cr: 0.01 to 2.00%,
Ni: 0.01 to 2.00%,
Cu: 0.01 to 2.00%,
Mo: 0.01 to 1.00%,
B: 0.0001 to 0.0100% and
W: 0.01 to 2.00%
High-strength steel sheet, characterized in that it contains one or two or more selected from the group consisting of. According to claim 1,
In mass%,
High-strength steel sheet, characterized in that it contains 0.0001 to 0.0100% of one or two or more selected from the group consisting of Ca, Ce, Mg, Zr, La and REM. According to claim 1,
A high-strength steel sheet characterized in that the austenite particles having an aspect ratio of 2.50 or less and a circle equivalent diameter of 0.80 µm or more have a density of 5.0 × 10 ¹⁰ particles / m 2 or less. A high-strength galvanized steel sheet, wherein a galvanized layer is formed on the surface of the high-strength steel sheet according to any one of claims 1 to 5. The method of claim 6,
High-strength galvanized steel sheet, characterized in that the Fe content in the galvanized layer is 3.0% by mass or less. The method of claim 6,
High-strength galvanized steel sheet, characterized in that the Fe content in the galvanized layer is 7.0% by mass or more and 13.0% by mass or less.